Process for making isotropic negative thermal expansion ceramics

ABSTRACT

Ceramic monoliths are described which exhibit tunable coefficients of thermal expansion from about −5 to −11×10 −6 ° C. −1  near ambient temperatures. These two-phase ceramics, which are fabricated, for example, by reactive sintering of WO 3  and ZrO 2 , consists of a matrix of ZrW 2 O 8  with inclusions of ZrO 2  having diameters less than 10 μm. Additives may increase the density of the monoliths to greater than 98% of the calculated density. Green body densities, pre-sintered particle size distribution, sintering atmosphere, microstructure, and mechanical properties are discussed. These ceramics may be used as substrates for thermally compensating fiber Bragg gratings.

This application is a divisional of application Ser. No. 09/146,518,filed Sep. 3, 1998, now U.S. Pat. No. 6,258,743.

FIELD OF THE INVENTION

This invention pertains to isotropic, negative thermal expansionceramics, and to a process for preparing isotropic, negative thermalexpansion ceramics which may, for instance, be used in temperaturecompensating grating packages.

TECHNOLOGY REVIEW

As described in U.S. Pat. No. 5,694,503, incorporated herein byreference, it is possible to package a fiber grating on or in a negativeexpansion material so that the package and grating dimensions decreasewith an increase in temperature, resulting in minimal variation of thereflection wavelength with temperature.

In order for this approach to be practical the value for the thermalexpansion of the package/substrate material must lie within anacceptable range. The ideal expansion coefficient is given by theexpression$\alpha_{ideal} = {\alpha_{fiber} + \frac{- A}{\left( {1 - P_{e}} \right)\quad \lambda_{nom}}}$

where α_(fiber) is the thermal expansion of the fiber grating (forinstance, 0.55×10⁻⁶° C.⁻¹), A is the temperature sensitivity of theunpackaged grating (e.g., 0.0115 nm °C.⁻¹ for a particular 1550 nmgrating), P_(e) is the photoelastic constant (typically 0.22) andλ_(nom) is the nominal grating wavelength (˜1550 nm in many cases). Fora particular 1550 nm grating an “ideal” package material would have athermal expansion coefficient (CTE) of −8.96×10⁻⁶° C.⁻¹. A packagematerial will still be beneficial even if its thermal expansion is notexactly equal to the ideal value. For the above assumptions a factor ofabout 20 improvement in temperature sensitivity would be achieved if thepackage material's thermal expansion coefficient were within 0.47×10⁻⁶°C.⁻¹ of the ideal value. Such improvement in thermal stability of afiber grating would be of commercial importance.

It is known that ZrW₂O₈ is metastable at room temperature, with thelower limit of stability being at 1105±3° C., below which ZrW₂O₈decomposes into ZrO₂ and WO₃. See, for instance, J. Graham et al., J.American Ceramics Society, Volume 42, page 570 (1959), and L. L. Y.Chang et al., J. American Ceramics Society, Volume 50, page 211 (1967).It is also known that ZrW₂O₈ has a relatively high and isotropicnegative coefficient of thermal expansion (CTE) over an extensive rangeof temperatures that includes room temperature. See, for instance, C.Martinek et al., J. American Ceramics Society, Volume 51, page 227(1968) and T. A. Mary et al., Science, Volume 272, page 90 (1996).Specifically, the CTE is substantially constant from near absolute zerotemperature to 150° C., with a value near −10×10⁻⁶° C.⁻¹. The materialexhibits an order-disorder transition at 150° C., after which the CTEdrops to −5×10⁻⁶° C.⁻¹. This value of the CTE is maintained until thedecomposition of ZrW₂O₈ which occurs at a relatively high rate near 800°C.

In view of this complex behavior of ZrW₂O₈, it is not surprising thatearlier attempts to produce mechanically strong monolithic bodies ofZrW₂O₈, that have a predetermined negative CTE, did not yield fullysatisfactory results. For instance, C. Verdon et al., ScriptaMaterialia, Volume 36, page 1075 (1997) report that their attempts toform electrically conducting bodies from ZrW₂O₈ and Cu with a low CTEresulted in decomposition of the ZrW₂O₈ and formation of Cu₂O along withother compounds. Such decomposition is generally undesirable and hindersthe production of suitable bodies.

To the best of our knowledge, prior art efforts to make ZrW₂O₈ ceramicbodies used the conventional technique of sintering ZrW₂O₈ powder. Thusproduced bodies have densities less than 90% of the theoretical densityand exhibit relatively poor mechanical properties, specifically, a lowmodulus of rupture. Additional prior art describes the preparation ofZrW₂O₈ powder from oxide precursors to be an incomplete reaction.

In view of the importance, for instance, a reduction of the temperaturedependence of the reflection wavelength of optical fiber gratings, itwould be highly desirable to be able to reproducibly make mechanicallystrong ceramic bodies having tunable negative CTE values, the ceramicbodies being useful, for example, for packaging of fiber gratings. Thisapplication discloses a method for making such bodies.

SUMMARY OF THE INVENTION

We have made the surprising discovery that reactive sintering ofappropriate percursor powders (e.g., ZrO₂ and WO₃) can result in(negative CTE) bodies (e.g., ZrW₂O₈) with substantially improvedproperties, as compared to analogous bodies produced by the prior artsintering techniques.

To the best of our knowledge, the prior art does not provide anysuggestion that the use of reactive sintering could provide the observedimproved results. By “reactive sintering” we mean herein compacting theunsintered body of the precursor oxides, rather than the powder of thedesired final phase, and then forming the desired phase and densifyingthe body in a single heat treatment step.

In a broad aspect of the invention is embodied in a method of makingnegative CTE ceramic bodies, and in bodies produced by the method.

More specifically, the invention is embodied in a method of making aceramic body having isotropic negative thermal expansion, the bodyhaving a major constituent selected from the group consisting of ZrW₂O₈,HfW₂O₈, ZrV₂O₇ and HfV₂O₇. The method comprises the steps of providing apowder mixture, forming a “green” body that comprises the powdermixture, and heat treating the green body. The powder mixture comprisesa first and a second oxide precursor powder, selected respectively fromthe group consisting of ZrO₂ powder and HfO₂ powder, and the groupconsisting of WO₃ powder and V₂O₅ powder. The heat treatment of thegreen body includes heating the body to a temperature below the meltingtemperature of the selected constituent, such that the selected majorconstituent is formed from the green body by reactive sintering.

By a “green” body we mean herein the compacted precursor powders as amonolithic body prior to sintering.

In a preferred embodiment the powder mixture has a non-stoichiometriccomposition (e.g., excess ZrO₂), and the heat treatment results information of a 2-phase material, e.g., ZrW₂O₈ majority phase, with ZrO₂inclusions dispersed in the majority phase. This embodiment allowstailoring of the CTE of the body.

In a further preferred embodiment the powder mixture comprises a minoramount of a sintering aid (e.g., Y₂O₃, Bi₂O₃, Al₂O₃, ZnO, TiO₂, SnO₂),whereby the density of the sintered body is substantially increased.

Important note for processing

Due to the unusually narrow stability region of ZrW₂O₈, standardtechniques for synthesizing and densifying composite ceramics containingZrW₂O₈ produce monoliths with inadequate physical properties. Highpurity ZrW₂O₈ is thermally stable between 1105 and 1260° C. Above 1260°C. ZrW₂O₈ peritectically decomposes into a Liquid phase, and below 1105°C. it decomposes into ZrO₂ and WO₃. Additives significantly alter thestability region of ZrW₂O₈; this study demonstrates a workingtemperature range of 1140 to 1180° C. for Y₂O₃ doped ZrW₂O₈.

BRIEF DESCRIPTION OF THE DRAWING

FIG. 1. Percent densities of ZrW₂O₈.xZrO₂ monoliths with Y₂O₃ additiveplotted with respect to weight percent of additional ZrO₂.

FIG. 2. SEM micrograph of a ZrW₂O₈.18 wt. % ZrO₂ monolith viewed withsecondary electron imaging. ZrO₂ inclusions with various particle sizeshaving a maximum diameter of approximately 10 μm appear with darkercontrast.

FIG. 3. Relative thermal expansion of a ZrW₂O₈.9.5 wt. % ZrO₂ monolithtaken over an extensive temperature range. The coefficient of thermalexpansion (CTE) for the range −100 to 100° C. and 200 to 300° C. is−10×10⁻⁶ and 3×10⁻⁶° C.⁻¹, respectively. A reversible order-disordertransition takes place near 140° C. which does not compromise themechanical strength of the monolith.

FIG. 4. Relative thermal expansion over ambient working temperatures forseveral ZrW₂O₈.xZrO₂ monoliths. The following compositions are presented(label, wt. % excess ZrO₂, volume fraction ZrO₂): a, 37.0%, 0.337; b,19.5%, 0.174; c, 9.5%, 0.084, d, 0%, 0. Nonlinearity of the thermalexpansion is enhanced as the zirconia content increases.

FIG. 5 illustrates the dependence of the Coefficient of ThermalExpansion (CTE) between 0 and 100° C. of two-phase ZrW₂O₈.xZrO₂ ceramicsas weight percent of additional ZrO₂. The change in CTE demonstrates alinear relationship to the relative amount of ZrO₂ inclusions.

FIG. 6. Reflection wavelength of a fiber Bragg grating as a function oftemperature is compensated by a ZrW₂O₈ 9.5 wt. % ZrO₂ substratethroughout ambient temperature range.

DETAILED DESCRIPTION OF THE INVENTION

The synthesis of ZrW₂O₈ has been described as challenging due itsmetastability. However, the fabrication of monoliths via reactivesintering is straightforward and reproducible upon consideration of thefollowing points: additives (including impurities), particle size, andsintering atmosphere. Similarly, reactive sintering may be used toprepare HfW₂O₈, ZrV₂O₇, and HfV₂O₇.

Additives

In phase equilibria, additives which form a eutectic may be utilized asa liquid phase sintering aid, thereby significantly increasing thedensity of a ceramic. Liquid phases accelerate material transport duringsintering by serving as a solvent for the solid phase with subsequentprecipitation to yield dense ceramics. Monoliths of ZrW₂O₈.xZrO₂prepared with a Y₂O₃ additive have achieved, densities greater than 99%the calculated theoretical density (FIG. 1). The density of a highpurity ZrW₂O₈ monolith is only 92%. Several oxides typically used assintering aids including refractory oxides were investigated (Table 1).Increased densities were observed without decomposition of ZrW₂O₈ forsmall amounts (0.1 wt. %) of additives. In order to easily observeenhanced differences in density, the samples were milled for only twohours (vide infra). Of the additives which were surveyed, all increasedthe density of the monoliths with exception of SiO₂. The followingseries illustrates the relative effectiveness of the additives:Y₂O₃>Bi₂O₃>Al₂O₃˜ZnO>TiO₂.

Densification via liquid phase sintering has been demonstrated inseveral ceramic systems and is extensively employed in the ceramicindustry in order to decrease sintering times and temperatures. Foreffective liquid phase sintering, only 1 vol. % of liquid within grainboundaries is necessary to aid densification. Increasing theconcentration of Al₂O₃, TiO₂, or Y₂O₃, to a few weight percent leads tocomplete decomposition of ZrW₂O₈ and melting of the monolith at thesintering temperature. This suggests that a liquid phase is present at1180° C. even at low concentrations of additives. Similar equilibriahave been observed in the alkali metal-WO₃—ZrO₂ systems with eutectictemperatures below 600° C. Consequently, liquid phase sintering is apreferred mechanism for enhanced densification.

Due to the impact of additives on the sintered density, it is importantto characterize the impurities in the starting materials and grindingmedia. An important consideration is the selection of appropriategrinding media based on composition in order to prevent undesiredimpurities, such as Al₂O₃ and calcium stabilized ZrO₂, fromincorporation into the monoliths.

Pre-sintering Particle Size

The particle size, particle size distribution, and uniformity ofparticle packing are important factors in the densification process. Inthis study of the preparation of ZrW₂O₈.xZrO₂ monoliths, these factorshave the greatest impact on sintered density. In general, smallerparticles are more reactive and tend to densify at lower temperatures.Coarse particles have less surface energy which diminishes the drivingforce for consolidation. For conventional powder processing, the optimumparticle diameter is approximately 1 μm. Particles significantly smallerthan this are difficult to pack uniformly and densely due toagglomeration and undergo excessive shrinkage with entrapped porosityduring firing. In addition, a range of particle sizes is advantageoustowards achieving maximum packing density in the green body. However,non-uniformities in particle packing may result in voids which aredifficult to eliminate during the densification process. Variations ingreen body densities produced during the forming process can also resultin warpage during sintering.

The dependence of sintered density on green body density wasaccomplished by dry-pressing samples at applied pressures of 1000 psi to20,000 psi. All the samples were pressed in the same ¼ in. cylindricaldie and therefore had the same green diameter. Information in Table 2expounds the effects of green body density on linear shrinkage and theaspect ratio of the pellets after sintering. Samples pressed at lowerpressures had lower green densities and underwent greater linearshrinkage during firing. All samples consolidated to a similar densityof approximately 4.45 g/cm³, independent of their initial greendensities. The initial particle size or surface area, rather thaninitial green density of the compact had a greater impact on the finalsintered density. These data suggest that the primary driving force fordensification of ZrW₂O₈ is the reduction of surface free energy.However, complexities due to reactive sintering with the aid ofadditives obscure the sole effect of surface energy minimization.

The as-received mixture of powders had a bimodal particle sizedistribution around 600 μm and 50 μm. Monoliths fabricated from thiscoarse, unreactive material actually decreased in density aftersintering. It was necessary to comminute the starting materials byvibratory milling. As the milling time increased the particle sizeinitially decreased significantly to approximately 1 μm diameter.Milling times were limited to 16 h due to concerns of contamination fromthe grinding media.

Sintering Atmosphere

Loss of oxygen in ZrW₂O₈ has been observed at temperatures as low as500° C. in vacuum, while retaining the crystal structure. Another studyfound that oxygen scavengers (e.g. Cu) promote decomposition of ZrW₂O₈at low temperatures (600° C.) through the formation of Cu₂O. It has beenfound that monoliths sintered in air at 1180° C. partially decompose onthe surface, followed by volatilization of WO₃. Heating ZrW₂O₈ in a N₂atmosphere at the sintering temperature leads to complete decompositioninto ZrO₂ and WO₃, whereas sintering in pure oxygen produced singlephase ZrW₂O₈ with minimal surface decomposition. In addition, monolithssintered in oxygen demonstrated reproducible physical properties.

Microstructure

The microstructure of a nearly stoichiometric ZrW₂O₈ monolith revealslarge grains of ZrW₂O₈ having approximately a 20 μm diameter which formthe matrix of the ceramic. The uniform distribution of pores and grainsizes is indicative of a homogeneously sintered material. A SEMmicrograph of a ZrW₂O₈.18 wt. % ZrO₂ monolith is presented in FIG. 2 inwhich excess zirconia is observed as nearly spherical inclusions havingvarious diameters. Individual ZrW₂O₈ grains cannot be distinguished inthe SEM image.

For the ZrW₂O₈.xZrO₂ compositions surveyed, the ZrO₂ volume fraction wasbelow the percolation limit such that a 0-3 connectivity (isolatedinclusions of ZrO₂ in a matrix of ZrW₂O₈) was preserved. The percolationlimit for a second phase with monodisperse diameters has been calculatedto exist at a critical volume fraction of 0.183. Although we haveprepared samples with a greater volume fraction of zirconia, allevidence indicates isolated inclusions. Thus significantly higher volumefractions are necessary to reach the percolation limit which is due tothe distribution of diameters of the inclusions. Beyond the percolationlimit the properties of a monolith with two interpenetrating3-dimensional matrices (3—3 connectivity) may demonstrate anomalies. Inaddition, it is possible that the stress placed on a 3—3 ceramic due tothe tetragonal to monoclinic phase transformation of the ZrO₂ matrixwould lead to catastrophic failure of the monoliths.

Powder X-ray diffraction indicates the presence of only the α—ZrW₂O₈phase (Powder Diffraction File, Joint Committee on Powder DiffractionStandards, JCPDS, Swarthmore, Pa., card number 13-557) and monoclinicBaddeleyite ZrO₂ phase (JCPDS card number 37-1484). However, smallamounts of additional phases may be present. The lower detectable limitof other phases by X-ray diffraction is estimated to be 1%.

Thermal Expansivity

ZrW₂O₈ has demonstrated a negative thermal expansion from 0.3 to 1050 K.The temperature dependence of the thermal expansion of a monolith ofZrW₂O₈ with 9.5 wt. % excess ZrO₂ is shown in FIG. 3. The order-disorderphase transition at 150° C. is reversible and does not compromise themechanical strength of the monoliths. In addition, cracking due to thelarge difference in thermal expansion of ZrW₂O₈ and ZrO₂ was notobserved over the temperature range of −100° to 300° C. The linearnegative thermal expansion of the composite facilitates its utilizationin applications.

The coefficient of thermal expansion of diphasic ceramic monoliths canbe tuned by compensating the large negative thermal expansion of ZrW₂O₈with a material having a positive CTE such as ZrO₂. Zirconia was chosenas the second phase for its thermodynamic stability in the presence ofZrW₂O₈ and for the ease of processing via the reactive sinteringtechnique. Several other refractories react with ZrW₂O₈ at the sinteringtemperature and therefore lead to irreproducible results due todecomposition. In addition to stability, the thermal expansion of ZrO₂is roughly linear from 20 to 100° C. with a CTE of 8×10⁻⁶° C.⁻¹. This isadvantageous since the compensation of the thermal expansion leads to anaveraged CTE which is also linear over ambient temperatures (FIG. 4).These diphasic ceramics can be tuned from −5 to −11×10⁻⁶° C.⁻¹ for densemonoliths. A linear correlation between the CTE and the volume fractionof ZrO₂ occurs over a broad range of compositions (FIG. 5). A simplemodel, based on an additive relationship of thermal expansions for amonolith without zirconia with a CTE of −11×10⁻⁶° C.⁻¹ and pure ZrO₂with a CTE of 8×10⁻⁶° C.⁻¹, would have a slope of 19. The linear fit toour data has a slope of 17.6 which is in close agreement for thecompositions surveyed. This approach, which neglects any anomalies atthe percolation limit, suggests that a monolith with a volume fractionof ZrO₂ of 0.58 would have a zero CTE.

The thermal expansion was measured along three directions normal to thefaces of the ceramic bars. Although ZrW₂O₈ crystallizes in a cubic spacegroup, the monoliths are prepared by uniaxially compressing thepre-sintered powder mixture. This processing could introduce anisotropicmacroscopic properties. However, the thermal expansion of the barsdemonstrates isotropic behavior. Heating rates up to 20° C./min totemperatures of 400° C. with a TMA analyzer did not reveal anydecomposition or cracking which would be distinguished asirreproducibility or discontinuities in the CTE, respectively.Degradation of the ceramics was not observed through several heating andcooling cycles.

Compensating the CTE with excess WO₃ has been accomplished, but themonoliths have extremely high porosity at the surface to a depth ofapproximately 120 μm. In order to fabricate mechanically robustmonoliths, extensive machining is necessary to remove the porous layer.

Fiber Grating Package

The Bragg wavelength (λ) in a vacuum is given by

λ=2n_(eff)Λ

where n_(eff) is the effective refractive index for the guided mode inthe fiber, and Λ is the period of the index modulations of the fiber(˜0.5 μm for a particular 1550 nm grating). The Bragg wavelength of afiber Bragg grating is temperature dependent primarily due to thetemperature dependence of the refractive index of the silica basedglass. In addition, the Bragg wavelength is strain dependent by alteringthe fringe spacing. As the temperature increases the refractive index ofglass increases and vice versa. Also, due to the positive thermalexpansion of silica (CTE˜0.5×10⁻⁶° C.⁻¹), the fringe spacing increasesslightly with temperature. The wavelength shift that corresponds torefractive index (n) changes due to temperature (T) variations andthermal expansion of silica glass (α_(thermal)) can be calculated asfollows: λ = 2n  Λ$\frac{\lambda}{T} = {{2\Lambda \quad \frac{n}{T}} + {2n\quad \frac{\Lambda}{T}}}$${\Delta\lambda} = {{2\Lambda \quad \frac{n}{n}\frac{n}{T}\Delta \quad T} + {2n\frac{\Lambda}{\Lambda}\frac{\Lambda}{T}\Delta \quad T}}$${\Delta\lambda} = {{\lambda \left( {{\frac{1}{n}\quad \frac{n}{T}} + \alpha_{thermal}} \right)}\Delta \quad T}$${{where}\quad {CTE}} = {\alpha_{thermal} = {\frac{1}{\Lambda}\frac{\Lambda}{T}}}$

For a 100° C. temperature change the shift in wavelength for aparticular 1550 nm grating is measured to be about 1.15 nm. The portionof this shift due to the CTE of the silica fiber (α_(thermal)) is ˜0.08nm, which is less than 8% of the total shift.

As mentioned before, the grating wavelength is also sensitive to strain.If a grating is stretched, then the grating wavelength will increase.The strain and wavelength relationship can be presented as follows:$\frac{\Delta \quad \lambda}{\lambda} = {\frac{\Delta \quad l}{l}\left( {1 - P_{e}} \right)}$

where Δl/l=ε, the strain, and P_(e) is the effective photo-elasticconstant (˜0.22) Due to this effect, it is possible to package a fibergrating on or in a negative expansion material so that package andgrating dimensions decrease with an increase in temperature, resultingin a value of Λ that falls as the temperature increases. By choosing theappropriate negative expansion coefficient (˜−9×10⁻⁶° C.⁻¹) thatmaintains a constant n_(eff)Λ product, a grating whose reflectingwavelength shows minimal variation with temperature can be achieved.

The dependence of the reflection wavelength of both a compensated and anuncompensated fiber Bragg grating as a function of temperature is shownin FIG. 6. The grating that is compensated by a ZrW₂O₈.9.5 wt. % ZrO₂monolith demonstrates a 0.05 nm deviation from −40 to 80° C., which isnearly ideal. This substrate can be utilized to fabricate a thermallycompensated package suitable for WDM (wavelength division multiplex)applications.

Mechanical Strength

Four-point bending tests were performed on bars of ZrW₂O₈. Preliminarytests revealed a low modulus of elasticity (0.5×10⁶ psi) for theZrW₂O₈.xZrO₂ monoliths. The modulus of rupture was determined to be 3000psi. With regards to the grating package, an applied force of 89 N isrequired to break the monolith which is much greater than the maximumtensile force which is exerted by the fiber (2 N).

The monoliths demonstrate brittle failure having an approximateconchoidal fracture surface. The mechanism proceeds via intergranularfracture around ZrO₂ inclusions and intragranular throughout the ZrW₂O₈matrix. Therefore shearing at boundaries between ZrW₂O₈ grains does notoccur at room temperature.

Conclusion

Reactive sintering of WO₃ and ZrO₂ powders produces dense monoliths withadequate strengths. The reactive sintering technique circumvents theinherent metastability of ZrW₂O₈ in the early stages of densification,thereby yielding reproducible fabrication conditions. Monolithscontaining zirconia inclusions demonstrate a range of thermal expansioncoefficients linearly related to the volume fraction of ZrO₂. A monolithof ZrW₂O₈.xZrO₂, which exhibits a negative thermal expansion in thedesired range, has been successfully prepared and shown to thermallycompensate a fiber Bragg grating.

TABLE 1 Effect of Sintering Aids on the Density of Monoliths Additive(0.1 wt. %) Density (g/cm³) None 3.25 SiO₂ 3.10 TiO₂ 3.82 ZnO 3.94 Al₂O₃3.95 Bi₂O₃ 4.05 Y₂O₃ 4.21

TABLE 2 Effect of Green Body Density on Shrinkage Applied SinteredLinear Pressure Diameter Shrinkage Aspect Density (psi) (mm) (%)^(a)Ratio^(b) (g/cm³)  1000 5.94 6.45 1.13 4.49  5000 6.02 5.20 1.17 4.4210000 6.08 4.25 1.24 4.51 20000 6.20 2.36 1.32 4.41 ^(a)Linear shrinkagedetermined along diameter. ^(b)Aspect ratio defined as diameter/heightof cylindrical monolith after sintering.

EXAMPLES

The following examples are presented to assist those skilled in thistechnology to understand and practice the invention, without in any wayintending to limit the invention to the exemplified embodiments.

Example 1

Preparation of ZrW₂O₈ composite ceramics via reactive sinteringtechnique utilizing methyl ethyl ketone (MEK) as the milling solvent.

Milling of ZrO₂ and WO₃ mixture in methyl ethyl ketone (MEK) for 10-20hours utilizing stabilized ZrO₂ grinding media.

Addition of approximately 2 wt. % of an organic binder (QPAC-40, PACPolymers Inc.) which is soluble in MEK to above mixture,

Evaporation of MEK during continuous stirring of mixture,

Sieving of dried mixture through 30-100 mesh screen,

Pressing of powder mixtures to form green body,

Firing in an oxygen atmosphere on a bed of coarse ZrW₂O₈ grains on Ptfoil,

Slowly heating said green body around 50° C. per hour to around 250° C.,

Subsequent heating of said green body around 500° C. per hour to around1150 to 1200° C. with an optimal sintering temperature of 1180° C.,

Holding said green body at 1180° C. for around 5 hours,

Cooling quickly to room temperature by withdrawing the sintered monolithfrom the furnace.

Example 2

Preparation of ZrW₂O₈ composite ceramics via reactive sinteringtechnique utilizing water as the milling solvent.

Milling of ZrO₂ and WO₃ mixutre in water for 10-20 hours utilizingstabilized ZrO₂ grinding media,

Addition of approximately 5-10 wt. % of an organic binder (polyvinylalcohol) which is soluble in water to mixture,

Evaporation of water during continuous stirring of mixture,

Sieving of dried mixture through 30-100 mesh screen,

Pressing of powder mixture to form green body,

Firing in an oxygen atmosphere on a bed of coarse ZrW₂O₈ grains on Ptfoil,

Slowly heating said green body around 50° C. per hour to around 250° C.,

Subsequent heating of said green body around 500° C. per hour to around1150 to 1200° C. with an optimal sintering temperature of 1180° C.,

Holding said green body at 1180° C. for around 5 hours,

Cooling quicly to room temperature by withdrawing the sintered monolithfrom the furnace.

Example 3

Extrusion of pre-reacted powders.

Milling of ZrO₂ and WO₃ mixture in methyl ethyl ketone (MEK) for 10-20hours utilizing stabilized ZrO₂ grinding media,

Evaporation of MEK,

Sieving of dried mixture through 30-100 mesh screen,

Addition of approximately 0.3 wt. % plasticizer and lubricant (UnionCarbide, PEG-400) to powder mixture.

Addition of approximately 1.0 wt. % dispersant (Angus Chemical Co.,AMP-95) to powder mixture,

Addition of approximately 4.5 wt. % binder (Rohm and Haas, B-1051) topowder mixture,

Addition of approximately 2.4 wt. % binder (Rohm and Haas, B-1052) topowder mixture,

Addition of approximately 6.1 wt. % water to powder mixture.

Blended under low shear conditions,

Extruded through die at room temperature.

Example 4

Preparation of ZrV₂O₇ composite ceramics via reactive sinteringtechnique utilizing water as the milling solvent.

Milling of ZrO₂ and V₂O₅ mixture in water for 10-20 hours utilizingstabilized ZrO₂ grinding media,

Addition of approximately 5-10 wt. % of an organic binder (polyvinylalcohol) which is soluble in water to mixture,

Evaporation of water during continuous stirring of mixture,

Sieving of dried mixture through 30-100 mesh screen,

Pressing of powder mixture to form green body,

Firing in an oxygen atmosphere on a bed of coarse ZrV₂O₇ grains on Ptfoil,

Slowly heating said green body around 50° C. per hour to around 250° C.,

Holding said green body at 850-900° C. for around 5 hours,

Cooling quickly to room temperature by withdrawing the sintered monolithfrom the furnace.

It is understood that various other modifications will be apparent toand can readily be made by those skilled in the art without departingfrom the scope and spirit of this invention. Accordingly, it is notintended that the scope of the claims appended hereto be limited to thedescription as set forth herein, but rather that the claims be construedas encompassing all the features of patentable novelty that reside inthe present invention, including all features that would be treated asequivalents thereof by those skilled in the art to which this inventionpertains.

What is claimed is:
 1. A method of making a ceramic body having anisotropic negative coefficient of thermal expansion, the body comprisinga first phase selected from the group consisting of ZrW₂O₈,ZrV₂O₇ andHfV₂O₇ and a second dispersed phase comprising at least one oxideselected from the group consisting of WO₃, V₂O₅, ZrO₂ and HfO₂,comprising the steps of, a) preparing a powder mixture comprising afirst and a second oxide precursor powder, the first oxide precursorpowder selected from the group consisting of ZrO₂ powder and HfO₂powder, and the second oxide precursor powder selected from the groupconsisting of WO₃ powder and V₂O₅ powder; b) forming a green bodycomprising said powder mixture; and c) heat treating the green body at atemperature below the melting temperature of the first phase, andforming the first phase from the heat treated green body by reactivesintering.
 2. The method according to claim 1, wherein the powdermixture contains a non-stoichiometric amount of one of the first andsecond precursor powders.
 3. The method according to claim 1, whereinthe powder mixture further comprises a sintering aid selected to yieldan increased density of the ceramic body.
 4. The method according toclaim 3, wherein the sintering aid is selected from the group consistingof Y₂O₃, Bi₂O₃, Al₂O₃, ZnO, TiO₂, SnO₂ and combinations thereof.
 5. Themethod according to claim 1, wherein said powder mixture is prepared byvibrational milling of said first and second oxide precursor powders ina milling solvent to form a slurry, adding a binder to said slurry, andevaporating said milling solvent to form a powder mixture.
 6. The methodaccording to claim 1, wherein said green body is formed by sieving saidpowder mixture to form a sieved powder, and pressing said sieved powderto form a green powder.
 7. The method according to claim 1, wherein saidheat treatment comprises firing said green body in an oxygen atmosphere,slow heating said fired body to about 250° C., subsequently heating toabout 1150° to 1200° and soaking said body for about 5 hours, andquenching said heated body to room temperature to form said ceramic bodyhaving an isotropic negative coefficient of thermal expansion.
 8. Themethod according to claim 1, wherein the first oxide precursor is ZrO₂.9. The method according to claim 1, wherein the first oxide percursor isHfO₂.
 10. The method according to claim 1, wherein the second oxidepercursor is V₂O₅.
 11. The method according to claim 1, wherein thesecond oxide percursor is WO₃.
 12. The method according to claim 1,wherein the amounts of first and second oxide precursors are selected toproduce a ceramic body having thermal expansion coefficient of about−9±2.5×10⁶° C.⁻¹.
 13. The method according to claim 1, wherein theamounts of first and second oxide precursors are selected to produce aceramic body having thermal expansion coefficient of about −9±1×10⁶°C.⁻¹.
 14. The method according to claim 1, wherein the amounts of firstand second oxide precursors are selected to produce a ceramic bodyhaving thermal expansion coefficient of about −9±0.25×10⁶° C.⁻¹.
 15. Themethod according to claim 1, wherein the first oxide is ZrO₂, the secondoxide is W₂O₃, and the sintering temperature is about 1180° C.
 16. Themethod according to claim 1, wherein the first oxide is ZrO₂, the secondoxide is V₂O₅ and the sintering temperature is about 850-900° C.